A typical prior art data storage system 10 used for longitudinal recording is illustrated in FIG. 1. In operation a magnetic transducer 20 is supported by a suspension (not shown) as it flies above a rotating magnetic disk 16. The magnetic transducer 20, usually called a “head” or “slider,” is composed of an element that performs the task of writing magnetic transitions (the write head 23) in ferromagnetic material on the magnetic disk, and another element that performs the task of reading the magnetic transitions (the read head 12) written in the ferromagnetic material on the magnetic disk. The magnetic transducer 20 is positioned by an actuator (not shown) over points at varying radial distances from the center of the magnetic disk 16 to read and write circular tracks (not shown). The magnetic disk 16 is attached to a spindle (not shown) driven by a spindle motor (not shown) to rotate the magnetic disk 16. The magnetic disk 16 comprises a substrate 26 on which a plurality of thin films 21 are deposited. The thin films 21 include ferromagnetic material in which the write head 23 writes the magnetic transitions, and in which the read head 12 reads the magnetic transitions.
There are three main categories of read heads 12, one current-in-plane (CIP) giant magnetoresistance (GMR) heads, one current-perpendicular-to-plane (CPP) GMR heads, and the other CPP tunneling magnetoresistance (TMR) heads. In each category, there are three types of read heads 12, one a top type, one a bottom type, and the other a dual type. A typical prior art top-type CIP GMR read head 12, as illustrated in FIG. 2, includes a bottom shield layer 38, a bottom gap layer 37, a top gap layer 41, and a top shield layer 39. Within the top and bottom read gap layers 37, 41, a top-type CIP GMR sensor 14 is located in a central read region, and hard-bias/lead layers 42, 43 are disposed in two side regions.
The top-type CIP GMR sensor 14, as shown in FIG. 2, comprises a nonmagnetic seed layer 31, a ferromagnetic free (sense) layer 32, an electrically conducting spacer layer 33, a ferromagnetic pinned (reference) layer 34, an antiferromagnetic pinning layer 35, and a nonmagnetic cap layer 36. GMR effects result from different magnetization orientations of the weakly coupled ferromagnetic free and pinned layers 32, 34 separated by the electrically conducting nonmagnetic spacer layer 33. The antiferromagnetic pinning layer 35 fixes the magnetization of the pinned layer 34 in a direction perpendicular to an air bearing surface (ABS) which is an exposed surface of the GMR sensor that faces the magnetic disk (the plane of the paper in FIG. 2). In a quiescent position when a sense current is conducted through the GMR sensor 14 without magnetic field signals from an adjacent rotating magnetic disk 16, the magnetization of the free layer is preferably parallel to the ABS. During GMR sensor operation with magnetic field signals from the adjacent rotating magnetic disk 16, the magnetization of the free layer is free to rotate in positive and negative directions from the quiescent position in response to positive and negative magnetic signal fields from the moving magnetic disk 16.
In the fabrication process of the top-type CIP GMR head 12, the top-type CIP GMR sensor 14 is deposited on the bottom gap layer 37 which is deposited on the bottom shield layer 38. The GMR sensor 14 typically comprises a Ta seed layer 31, Ni—Fe/Co—Fe ferromagnetic free layers 32, a Cu spacer layer 33, a Co—Fe pinned layer 34, an antiferromagnetic Ir—Mn, Pt—Mn or Ni—Mn pinning layer 35, and a Ta cap layer 36.
Photolithographic patterning and ion milling are applied to define the read region of the GMR sensor 14. The hard-bias/leads layers 42 and 43 are then deposited in the two side regions of the GMR sensor. The hard-bias/lead layers 42, 43 preferably comprise a Cr film, a ferromagnetic Co—Pt—Cr film, a Cr film, a Rh film, and a Ta film. An electrically insulating nonmagnetic top gap layer 41 is deposited over the cap layer 36 and hard bias/leads layers 42, 43. A top shield layer 39 is formed over the nonmagnetic top gap layer 41.
In this top-type CIP GMR sensor, ferromagnetic/antiferromagnetic coupling occurs between the pinned and pinning layers, producing a unidirectional anisotropy field (HUA). This HUA must be high enough to rigidly pin the magnetization of the pinned layer (M2) in a transverse direction perpendicular to an air bearing surface (ABS) for proper sensor operation. Ferromagnetic/ferromagnetic coupling also occurs across the spacer layer, producing a ferromagnetic coupling field (HF). This HF must be precisely controlled so that the sum of HF and a current-induced field (HI) counterbalances a demagnetizing field (HD) in the sense layer (HF+HI=HD), thereby orienting the magnetization of the sense layers (M1) in a longitudinal direction parallel to the ABS for optimally biased sensor operation. In a quiescent state, this GMR sensor exhibits a resistance of RO+RA, +(1/2)RG, where RO is a nonmagnetic resistance, RA is the maximum anisotropy magnetoresistance (AMR) of the free layers, and RG is the maximum giant magnetoresistance (GMR). When receiving a signal field from a magnetic disk, M1 rotates while M2 remains unchanged. This M1 rotation changes the resistance of the GMR sensor by ±ΔRG sin θ1−ΔRA sin2θ1, where θ1 is the angle of M1 rotation from the longitudinal direction.
When the GMR sensor is operating at elevated temperatures in the data storage system, an inadequate exchange coupling can cause canting of the magnetization of the pinned layer from the preferred transverse direction, causing malfunction of the sensor operation. The operation temperature of the GMR sensor in the data storage system can reach 180 degrees C. or more. A high HUA at high temperatures ensures proper sensor operation at high temperatures. This thermal stability is typically described by a blocking temperature (TB), where the ferromagnetic/antiferromagnetic exchange coupling diminishes and HUA is zero. A higher TB typically indicates a higher HUA at the sensor operation temperature.
The effort to increase the GMR coefficient, HUA and TB is typically directed to the selection of ferromagnetic and antiferromagnetic films from various alloy systems as pinned and pinning layers. Recently, a ferromagnetic 90Co-10Fe alloy film (in atomic percent) has replaced a ferromagnetic Co film as the preferred pinned layer, in order to increase the GMR coefficient, HUA and TB. An antiferromagnetic film selected from a Pt—Mn or Ni—Mn alloy system as a pinning layer has been extensively used in the GMR sensor.
In the selection process of an antiferromagnetic film from the Pt—Mn or Ni—Mn alloy system as a pinning layer, the Mn content of the Pt—Mn or Ni—Mn film must be carefully selected. A small difference in the Mn content leads to substantial variations in both HUA and TB. In addition, since the Mn is the most diffusive and corrosive chemical element among all the chemical elements used in the GMR sensor, its content substantially determines the corrosion resistance and thermal stability of the GMR sensor.
The currently used Mn content of the Pt—Mn or Ni—Mn films is selected only from a small composition range for attaining a high HUA. This Mn content may not be low enough to minimize the Mn diffusion, attain a high TB, and ensure high corrosion resistance. Hence, it is difficult, or almost impossible, to find a suitable Mn content for either the Co—Fe/Pt—Mn or Co—Fe/Ni—Mn films to attain a high HUA and a high TB simultaneously, as well as desirable corrosion resistance.
For example, in the prior art GMR sensor with a Ni—Mn pinning layer, a Mn content of more than 57 at % is selected in order to attain a high HUA beyond 600 Oe. However, previous studies indicate that such a high Mn content leads to a low TB and to a low corrosion resistance. Hence, to operate a smaller GMR sensor properly at high temperatures for magnetic recording at ever increasing densities, very robust pinning layers must be found.
In previous studies, a GMR sensor with a pinning layer formed of two antiferromagnetic films selected from two different binary alloy systems, such as Ir—Mn/Pt—Mn, Ir—Mn/Ni—Mn, Pt—Mn/Ni—Mn or Ni—Mn/Pt—Mn films, has been explored. The Ir—Mn film is selected and preferred to be in contact with the Co—Fe film since it does not require annealing for developing exchange coupling to the Co—Fe film, thereby eliminating concerns on the Mn diffusion. The Pt—Mn film is also selected to be in contact with the Co—Fe film to minimize the Mn diffusion and to provide a high HUA. The Ni—Mn film should not to be in contact with the Co—Fe film, but is preferably used to provide a high TB. However, since the Ir—Mn, Pt—Mn, and Ni—Mn films have different lattice parameters, the lattice mismatch causes exchange decoupling between the two different antiferromagnetic films, leading to difficulties in achieving the desired improvements.
A GMR sensor with a pinning layer of an antiferromagnetic film selected from a ternary alloy system, such as Ir—Pt—Mn, Ir—Ni—Mn, Pt—Ni—Mn films, etc., has also been explored. The antiferromagnetism has been found to be very weak, possibly due to incompatibility of Ir, Pt and Ni elements.
A Ni—Cr—Fe seed layer wherein the atomic percentage of Cr is between 20 and 50% is described by Lee, et al. in U.S. Pat. No. 6,141,191. A top-type GMR sensor is described with the structure of seed/free/spacer/pinned/AFM/cap layers, where the seed layer is a non-magnetic Ni—Cr—Fe or Ni—Cr film and the AFM layer is preferably a Ni—Mn film. The nonmagnetic Ni—Cr—Fe seed layer is said to result in a large-grain structure in the deposited layers enhancing the GMR coefficients and the thermal stability. The improved thermal stability enables the use of the Ni—Mn film with its high blocking temperature and strong pinning field as the AFM pinning layer, without performance degradation from the high temperature anneal step needed to develop the desired exchange coupling.
Huai, et al. (U.S. Pat. No. 6,222,707) disclose bottom-type and dual-type GMR sensors using Ni—Cr—Fe seed layers with a range of Cr atom percentage between 20% and 50%, with approximately 25 atomic percent being preferred. The top-type GMR sensor comprises seed/AFM/pinned/spacer/free/cap layers, while the dual-type GMR sensor comprises seed/AFM/pinned/spacer/free/spacer/pinned/AFM/cap layers. An improvement in the texture of synthetic pinned layers comprising Co—Fe/Ru/Co—Fe films is noted.
In U.S. Pat. No. 6,046,892 to Aoshima, et al. a bottom-type GMR sensor is disclosed with Co—Fe—B free and pinned layers and Ta/Ni—Cr—Fe seed layers.
In a published U.S. patent application Ser. No. 2004/0105193 by Horng, et al., a 30 angstrom thick seed layer of Ni—Cr with 31% atomic percent of Cr is used to form a bottom-type GMR sensor and a dual-type GMR sensor having synthetic pinned layers. The seed layer is said to allow the use of extremely thin (approximately 80 angstroms) Pt—Mn pinning layers as well as extremely thin pinned and free layers.
In a published U.S. patent application Ser. No. 2004/0042130 by Lin, et al. three seed layers comprising Al—O(3 nm), Ni—Cr—Fe(3 nm) and Ni—Fe(1 nm) films are followed by the Pt—Mn pinning layer. The '130 application is commonly assigned with the present application and has a common co-inventor with the present application. The Al2O3 film used as the bottom gap layer is preferably directly sputtered in an argon gas from an alumina target, while the Al—O film used as the seed layer is preferably reactively sputtered in mixed argon and oxygen gases from an aluminum target. A pinning layer, preferably comprising a 15 nm thick Pt—Mn film, is then deposited on the seed layers. Thereafter, pinned layers are deposited on the pinning layer. The pinned layers comprise a ferromagnetic Co—Fe first pinned layer, an antiparallel (AP) Ru spacer layer, and a ferromagnetic Co—Fe second pinned layer. A spacer layer, preferably a Cu—O film, is deposited on the second pinned layer. Thereafter, free layers, preferably comprising Co—Fe and Ni—Fe films, are deposited on the Cu—O spacer layer. The cap layers, preferably comprising Cu and Ta films, are then deposited on the free layers.
In order for the GMR sensor to attain a narrower read gap for a higher linear density, a 7.5 nm thick Ir—Mn pinning layer has been suggested to replace the 15 nm thick Pt—Mn pinning layer. However, the use of the Ir—Mn pinning layer appears to be impractical due to the poor corrosion resistance of the Ir—Mn pinning layer.
Ir—Mn, Pt—Mn and Ni—Mn pinning layers have been extensively used in the prior art, but some difficulties still remains in using these pinning layers for magnetic recording at ever increasing densities. FIG. 3 shows easy-axis high-field magnetoresistance (MR) responses of top-type GMR sensors comprising Ta(3)/Ni—Fe(4.5)/Co—Fe(0.6)/Cu(2.4)/Co(3.2)/Ir—Mn(7.5)/Ta(6), Ta(3)/Ni—Fe(4.5)/Co—Fe(0.6)/Cu(2.4)/Co(3.2)/Pt—Mn(25)/Ta(6) and Ta(3)/Ni—Fe(4.5)/Co—Fe(0.6)/Cu(2.4)/Co(3.2)/Ni—Mn(25)/Ta(6) films (thickness in nm) after annealing for 2 hours at 280° C. with a magnetic field of 800 Oe in a high vacuum oven. This anneal is not needed for the Ir—Mn pinning layer which contains an antiferromagnetic face-centered-cubic (fcc) phase after deposition, but must be used for the Pt—Mn and Ni—Mn pinning layers to cause a phase transformation from a nonmagnetic fcc phase to an antiferromagnetic face-centered-cubic (fct) phase. The Ir—Mn GMR sensor exhibits a very low HUA but the highest GMR coefficient. The Pt—Mn GMR sensor exhibits the highest HUA but a lower GMR coefficient. The Ni—Mn GMR sensor exhibits the lowest GMR coefficient and the lowest HUA.
FIGS. 4 and 5 show HUA and the GMR coefficient versus anneal time at 280° C. for the Ir—Mn, Pt—Mn and Ni—Mn GMR sensors. The temperature is controlled by a reactive heater and a continuous nitrogen flow in a tube oven attached to a vibrating sample magnetometer (VSM). In each anneal cycle, the oven temperature increases from ˜30° C. to 280° C. in two minutes, remains at 280° C. for a set time, and then decreases to 30° C. in two minutes. After each anneal cycle, GMR properties are measured at 30° C. After annealing for 5 minutes, the HUA of the Ir—Mn GMR sensor reaches its saturation value of 110 Oe, corresponding to JK (an intrinsic exchange coupling energy) of 0.04 erg/cm2. After annealing for 2 hours, the HUA of the Pt—Mn GMR sensor reaches its saturation value of 410 Oe, corresponding to JK of 0.16 erg/cm2. After annealing for 13.2 hours, the HUA of the Ni—Mn GMR sensor reaches as high as 620 Oe, corresponding to JK of 0.24 erg/cm2, without saturation yet. However, this extended anneal causes a substantial decrease in the GMR coefficient.
FIG. 6 shows HUA versus temperature for the Ir—Mn, Pt—Mn and Ni—Mn GMR sensors. The Ir—Mn, Pt—Mn and Ni—Mn GMR sensors exhibit TB (a blocking temperature, which HUA reaches 0) of 270, 370 and 380° C., respectively. When the Ir—Mn, Pt—Mn and Ni—Mn GMR sensors operate at 160° C., the HUA values are 60, 340 and 540 Oe, respectively. Hence, the Ir—Mn GMR senor shows the worst thermal stability, while the Ni—Mn sensor the highest. In addition, the Ir—Mn GMR sensor has poor corrosion resistance.
Based on all these experimental results, the Pt—Mn pinning layer appears to be the most suitable among the three types of pinning layers. As recently described in the prior art, the top-type Pt—Mn GMR sensor has been improved with four major approaches. First, the Ta seed layer has been replaced by Al—O/Ni—Cr—Fe/Ni—Fe seed layers for increasing the GMR coefficient. Second, the Co—Fe pinned layer is replaced by Co—Fe/Ru/Co—Fe pinned layers for minimizing a demagnetizing field (HD). Third, the Cu spacer layer has been replaced by a Cu—O spacer layer for attaining a negative ferromagnetic coupling field (HF). Fourth, the structure of the GMR sensor is reversed for ease in a read-head fabrication process. In this bottom-type GMR sensor, HD is smaller than HI, so that HF must be negative in order for the sum of HF and HD to counterbalance HI (HF+HD=HI) for optimally biased sensor operation.
The Ir—Mn pinning layer appears to be the least suitable among the three types of pinning layers. It nevertheless has two unique features important for magnetic recording at ever increasing densities. First, the Ir—Mn pinning layer can be as thin as 7.5 nm, so that it can be sandwiched into a narrow read gap for a high linear density. In contrast, the Pt—Mn and Ni—Mn pinning layers must be at least as thick as 15 and 20 nm, respectively. Second, the Ir—Mn pinning layer causes the Co—Fe pinned layer to induce an easy-axis coercivity (HCE) much lower than the Pt—Mn and Ni—Mn pinning layers. This low HCE is important in preventing the magnetization of the Co—Fe pinned layer from irreversible rotation.
A further improvement of the Ir—Mn pinning layer is thus desired for magnetic recording at ever increasing densities. In this invention, a Cr element is added into the Ir—Mn pinning layer for improving its corrosion resistance. On Al—O/Ni—Cr—Fe/Ni—Fe seed layers with optimal compositions and thicknesses, the Ir—Mn—Cr pinning layer as thin as 5 nm can strongly exchange-couple to the Co—Fe pinned layer with an optimal composition, inducing a high HUA.